Method for producing high speed steel

ABSTRACT

A method for producing a high speed steel that with reference to its chemical composition consists of the following elements: 1-3 wt-% carbon (C), 3-6 wt-% chromium (Cr), 0-7 wt-% molybdenum (Mo), 0-15 wt-% tungsten (W), 3-14 wt-% vanadium (V), 0-10 wt-% cobalt (Co), 0-3 wt-% niobium (Nb), 0-0.5 wt-% nitrogen (N), 0.2-1 wt-% yttrium (Y), and remainder iron (Fe) and unavoidable impurities, and wherein Mo+0.5W=2-10 weight %, characterized in that the method comprises the steps of: providing a powder comprising the elements of the high speed steel, forming a body of the powder, and subjecting the body to elevated heat and pressure such that a consolidation of the powder thereof is achieved.

TECHNICAL FIELD

This invention refers to a method for producing a high speed steel with the composition according to the preamble of claim 1.

BACKGROUND

New materials suitable for applications involving elevated temperatures and wear are needed today. Such application areas may, for example, include, hot forging tools for metal forming, combustion engine parts etc.

Several different special alloys suitable for high temperature use exist, for example FeCrAl-alloys, NiCrAl-alloys, Ni-base alloys, Co-base alloys and special stainless steels. However, the FeCrAl-, NiCrAl- and Ni-alloys are too soft to be used in the aforementioned areas of applications. Some Co-base alloys are sufficiently hard but too expensive to be a practical alternative for most applications.

High speed steel (HSS) offers good hardness at room temperature and is able to maintain that hardness up to 600° C. However, for some applications, it is desirable to maintain the room temperature hardness at temperatures well above 600° C.

In the art, it is known that alloying the high speed steel with cobalt improves its hardness at higher temperatures.

Several different ways to improve high speed steel usability in high temperature applications exists in the art.

It is known in the art that adding alloying elements such as cobalt (Co) in combination with strong carbide-forming elements which act through the base mass or matrix increase the high speed steel's hardness at high-temperature and thereby the wear resistance.

It is also known in the art that adding alloying elements such as tungsten, molybdenum and vanadium increases the high speed steel's ability to withstand high-temperature, i.e. their hot hardness and high temperature wear, due to formation of carbides based on said alloying elements.

It is also known in the art that an increase of the formed amount of carbides by alloying with carbide forming elements such as chromium, molybdenum, tungsten and vanadium positively contributes to the wear resistance of the high speed steel.

A large amount of alloying elements makes the high speed steel difficult to produce by standard casting techniques. The resulting microstructure suffers problems from severe segregation with very coarse carbides which results in very poor toughness and strength of the high speed steel. In applications, in which the steel is subjected to forging after the casting thereof, some of these problems are overcome as a result of the effect that such deformation has on the microstructure of the material.

For alloys, that in addition to high temperature and wear need to withstand an oxidative/corrosive environment, some new demands arise due to synergy effects. For materials that are simultaneously exposed to corrosion/oxidation and wear, the behaviour of the material is of utter importance. The oxidation kinetics as well as oxide scale mechanical properties and adhesion become important factors for the use of the alloy.

By depositing special coatings on the high speed steel, it's oxidation behaviour can be altered (Reduction of wear in critical engine components using ion-beam-assisted deposition and ion implantation, J. H. Arps et al., Surface and Coatings Technology 84 1996 p 579-583). However, such coatings are of limited use in applications with heavy wear, since they only bring a thin protective coating (in the micrometer range) to the steel. In applications in which the wear can be expected to be in the millimetre range, steel products with this type of coating will require subsequent, further coating as the initial coating would wear down. This would increase the cost of the product considerably over its expected lifetime.

Another feasible solution would be to utilize surface modification of the high speed steel, such as ion-implantation (Ion beam modification of metals, G. Dearnley, Nuclear Instruments and Methods in Physics Research B50 1990 p358-367). However, a problem with ion-implantation is the Gaussian distribution of ions, which causes depth varying material characteristics. Ion-implantation also implies a limited useful thickness of the modified layer; therefore it is not suited for use in the aforementioned applications.

U.S. Pat. No. 5,989,491 to Isamoto et al. discloses a method that uses oxide dispersion strengthening for a powder metallurgy alloy. The inventors behind this patent noted that fine dispersion of particles of an oxide in an oxide dispersion strengthened heat resisting powder metallurgy alloy enhanced creep rupture strength in addition to fundamental heat resisting properties inherent to heat resisting alloys. However, the alloy disclosed in U.S. Pat. No. 5,989,491 is not suitable for mechanical applications involving wear such as those aforementioned, since the wear resistance of the end product will not be affected by the addition of fine particles of an oxide.

Several patents disclose the use of rare-earth elements in high speed steels in connection with applications at elevated temperatures, for example see: JP1142055A, JP63213641A, CN101037760A, JP1159349A, JP1142056A, CN101078090A, JP6299298A, CN1693527A, JP57143468A, JP2005281839A, CN101831590A, JP8041592A, JP57143471A, JP2003253396A, and JP1008252A. However, all of these alloys are cumbersome to produce with standard casting techniques. The level of alloying some of these high speed steels exhibits, likely causes problems with segregation of alloying elements during solidification and form coarse carbide structures. Therefore, it is cumbersome to manufacture a high speed steel, with high levels of alloying elements and rare-earth elements, and simultaneously achieve uniform material properties and a well controlled microstructure.

JP57085952A (abstract) discloses an alloy with a composition corresponding to the composition according to the preamble of claim 1 of the present invention. It must be assumed that the document in question also discloses casting as the method of producing the alloy. It must be assumed that the steel disclosed therein has a microstructure that results in poor strength of the material and therefore makes it less usable as a wear part.

THE OBJECT OF THE INVENTION

The object of this invention is to present a method by means of which the above mentioned problems associated with manufacturing of high speed steel comprising of rare earth element yttrium (Y) for applications that involve wear at elevated temperatures, are reduced or solved.

The present invention also aims at providing a manufacturing method which increases the ability of high speed steel to withstand wear at elevated temperatures.

Thus, the present invention is based on the insight of the problems of segregation and coarse carbide structures associated with casting and the addition of yttrium into conventional high speed steel.

SUMMARY

The object of the invention is achieved by a method for producing a high speed steel that with reference to its chemical composition consists of the following elements, in weight %: Carbon (C) 1-3, Chromium (Cr) 3-6, Molybdenum (Mo) 0-7, Tungsten (W) 0-15, Vanadium (V) 3-14, Cobalt (Co) 0-10, Niobium (Nb) 0-3, Nitrogen (N) 0-0.5, Yttrium (Y) 0.2-1, and remainder consisting of iron (Fe) and unavoidable impurities and wherein Mo+0.5W=2-10, said method being characterised in that the method comprises the steps of: providing a powder comprising the elements of said high speed steel, forming a body of said powder, and subjecting said body to elevated heat (temperature) and elevated pressure such that a consolidation of the powder thereof is achieved. This step may be referred to as the consolidation step or a hot isostatic pressure (HIP) step.

During said consolidation step, the steel is in solid state, i.e. non-molten state. Preferably, the temperature during said step of elevated temperature is within the range of from 950-1200° C., wherein lower temperatures may be required for alloys with relatively high content of C and low content of alloying elements such as Mo, W, Co, Y, etc and wherein higher temperatures within said range is required for alloys with relatively low content of C and high content of said further alloying elements. If the temperature is too low, the final result will be a porous material, and if the temperature is too high, the material might start to melt, which should be avoided.

The pressure during the consolidation step is dependent on the temperature, which is chosen for each respective steel composition. A relatively low temperature may be compensated for by means of a higher pressure. Preferably, for the compositions within the scope of the present invention, as defined in claim 1 and also in the dependent claims, and for the mentioned temperature range of the consolidation step, the pressure should be in the range of from 800-1500 bar. In general, higher content of alloying elements will require a higher pressure for a specific chosen temperature.

Reference throughout the specification to “one embodiment” or “an embodiment” means that a particular feature, structure, or characteristic described in connection with an embodiment is included in at least one embodiment of the subject matter disclosed. Thus, the appearance of the phrases “in one embodiment” or “an embodiment” in various places throughout the specification is not necessarily referring to the same embodiment. Further, particular features, structures or characteristics may be combined in any suitable manner in one or more embodiments.

According to an embodiment, the body of consolidated powder, that now has a very low porosity level or no porosity at all, is then subjected to soft annealing. The soft annealing is performed in order to facilitate subsequent machining of the alloy. Preferably, the maximum temperature of the soft annealing step is the temperature of the foregoing consolidation step, while the minimum temperature is the temperature at which the steel undergoes softening and carbides in the steel spheroidize and the martensite transforms to ferrite. In any case, the temperature must not be so high that it results in severe coarsening of the carbide grain size.

The selected soft annealing temperature will depend on the composition of the alloy. Generally, higher contents of alloying elements will require a higher annealing temperature. Accordingly, for the compositions within the scope of the present invention, the soft annealing temperature will preferably be in the range of from 600 to 900° C. The duration of the soft annealing should be sufficiently long to reach sufficiently high ferrite content in the material. Preferably, after soft annealing the ferrite-austenite ratio should be at least 95/5. The steel is cooled relatively slowly, in order to avoid formation of martensite or bainite in the alloy. Preferably, the cooling rate is within the range of from 5-20° C./hour, depending on the composition of the alloy. Cooling with this rate is performed down to a temperature below, which the cooling rate will no longer affect the formation of bainite, martensite. Below that temperature, the cooling may be natural, and the cooling rate may depend only on the outer conditions reigning. For the alloys within the scope of the invention, this temperature may be in the range of from 600-700° C.

The body may thereafter be subjected to machining if necessary and thereafter heat treated with a hardening (austenizing) step at a temperature in the range of from 950-1200° C., depending on the specific composition of the steel that is hardened. After hardening, there will be some remaining austenite in the steel, the main part of the steel now being martensite. This austenite is removed by means of subsequent annealing steps. During the first step remaining austenite is transformed into martensite. However, this martensite being very brittle, this will require a further annealing step in order to become sufficiently ductile.

Depending on the composition and the amount of austenite that remains in the steel after hardening, the number and duration of the annealing steps may vary. According to a one embodiment of the invention, annealing steps are performed until the level of remaining austenite is maximum 5%, preferably maximum 2%.

The technical effect of the method of the present invention as disclosed hereinabove or hereinafter is that the rare earth element yttrium is evenly distributed in the powder. If the high speed steel according to the inventive concept would have been produced by a conventional casting method, the highly reactive element yttrium would segregate and not be evenly distributed. An even distribution of yttrium in the high speed steel base-matrix causes an oxide scale that is formed to adhere effectively to the high speed steel. The added yttrium also changes the growth kinetics of the oxide scale so that the scale quickly grows to a saturation thickness. The growth rate of the oxide scale is drastically reduced above this saturation thickness. The beneficial technical effect on the wear resistance, at elevated temperatures, due to the fine dispersion of yttrium in the base-matrix of the high speed steel is unexpectedly good. This technical effect is beyond what a person skilled in the art would expect from an addition of yttrium using a powder metallurgy method. In fact, the gain in technical effect is so high that it, unexpectedly, compensates for the higher costs related to the use of powder metallurgy as the method of producing this steel, making the steel very useful in any application in which it is subjected to severe wear conditions. In particular, the steel will have a mean carbide particle size which is much lower than that of a corresponding material made using casting method. According to the present invention, the steel should have a mean carbide particle size of <3 μm, something it will have if the method of the invention is being used for the production thereof. As a result of the production method of the present invention, the steel will also have an isotropic microstructure, which is also advantageous for its wear properties. In other words, the present invention teaches that the consolidation step and the subsequent heat treating steps shall be performed such that the steel obtains a mean carbide particle size which is <3 μm and an isotropic microstructure.

The properties of the formed oxide scale are extremely important in applications that besides high temperature and wear also include oxidation/corrosion. In oxidative/corrosive applications, it is of great importance that damages in the oxide scale are quickly repaired by a fast growth of the oxide scale itself, and this is achievable by using the material produced by the inventive method.

According to one embodiment of the inventive method, the provision of the powder mixture comprises the step of argon-atomisation of molten metal comprising said elements into said powder. By using argon-atomisation of the molten metal the amount of nitrides is minimized compared to using nitrogen-atomisation wherein the use of nitrogen gas causes the nitrides to form.

According to the present invention, the yttrium content of the high speed steel is within the range 0.20 to 1.0 weight %. It is preferred that the yttrium content of the high speed steel is more than 0.40 weight %, and less than 0.70 weight % more preferably less than 0.60 weight %. In one preferred embodiment, the yttrium content is with the range of from 0.45-0.60 weight %, such as from 0.4-0.5 weight %, such as 0.4, 0.41, 0.42, 0.43, 0.44, 0.45, 0.46, 0.47, 0.48, 0.50 weight %.

The yttrium content defined in the interval above gives the aforementioned positive effects on the oxide scale. Especially the yttrium content in the range of from 0.45-0.60 weight % gives a very good increase in the ability of the high speed steel to withstand high temperature wear. The lower limit 0.20% of the interval defines a starting point from where a significant positive effect of yttrium on the high temperature wear can be identified. The higher limit of 1% indicates the end of the interval from where a significant positive effect of yttrium on the high temperature wear can be identified.

According to an embodiment of the inventive method, the carbon (C) content of said high speed steel is in the range of from 1.1-1.4 weight %. The amount of carbon should be sufficient to form the carbides necessary for the wear resistance of the high speed steel. Preferably the amount of carbon should be enough to produce a high speed steel with sufficient hardenability. The lower limit of 1.1% defines a minimum carbon content in order to form a high speed steel with the desired carbides and hardenability. The higher limit of 1.4% defines maximum carbon content in this embodiment, above which austenite may be formed.

According to an embodiment of the inventive method, the chromium (Cr) content is in the range of from 3.0-6.0 weight %. This interval causes good hardenability as well as the necessary forming of carbides. However, too much chromium causes formation of residual austenite and increased risk for over-tempering; therefore the upper limit of Cr must not be exceeded. According to one embodiment, the Cr content is within the range of from 4.0-5.0 weight %.

According to an embodiment of the inventive method, the molybdenum (Mo) content is in the range of from 4.5-5.5 weight %. This interval causes secondary hardening by precipitation of carbides that will increase the hot hardness and wear resistance of the high speed steel.

According to an embodiment of the inventive method, the tungsten (W) content is in the range of from 6.0-7.0 weight %. This interval causes secondary hardening by precipitation of carbides that will increase the hot hardness and wear resistance of the high speed steel.

It is a well-known fact that Mo and W have similar effects on this kind of steel and that they are therefore to a large extent replaceable with each other. According to claim 1, Mo+0.5W=2-10 weight %. According to a preferred embodiment, Mo+0.5W=5-8.5 weight%. It should be pointed out that the elements having a lower limit of 0 weight % are optional.

According to an embodiment of the inventive method, the vanadium (V) content is in the range of from 3.0-5.0 weight %. This interval causes secondary hardening by precipitation of carbides that will increase the hot hardness and wear resistance of the high speed steel. However, too much vanadium causes the high speed steel to become brittle and therefore, the upper limit must not be exceeded. According to a preferred embodiment the V content is in the range of from 3.0-3.5 weight %.

According to an embodiment of the inventive method, the cobalt (Co) content of said high speed steel is in the range of from 8.0-9.0 weight %. The alloying of high speed steel with cobalt improves the tempering resistance and hot hardness, both of which are of great importance for the high speed steel to be used in a high temperature wear application. The amount of cobalt also has an effect on the hardness of the high speed steel by affecting the amount of retained austenite, causing said retained austenite to easily be converted to martensite during tempering. The selected interval for cobalt is a suitable interval for a high speed steel of this composition wherein the upper level is more an economic compromise than a scientific constraint. Alternatively, if cobalt is not to be used in the above-defined range, the cobalt content is 0% or at an impurity level.

Powder metallurgical high speed steel produced by the inventive method possesses properties, such as very good resistance to high temperature wear even in oxidative/corrosive environments.

BRIEF DESCRIPTIONS OF THE DRAWINGS

The inventive concept will now be further explained using reference figures in connection with attached drawings and graphs, in which

FIG. 1 is a schematic figure of a “pin on disc” test equipment,

FIG. 2 shows a cross section of a typical groove obtained from a “pin on disc” evaluation, perpendicular to the longitudinal direction,

FIG. 3 is a diagram showing the groove depth at room temperature and 650° C. for the alloys A, B and C in the “pin on disc” experiment,

FIG. 4 is a diagram showing the volume loss per meter at 650° C. for the alloys A, B and C in the “pin on disc” experiment, and

FIG. 5 shows the hardness in HRC for alloy A, B and C.

DETAILED DESCRIPTION

The industrial production of semi-finished products, components and cutting tools based on powder metallurgical high speed steel started 35 years ago. The first powder metallurgical production of high speed steel was based on hot isostatic pressing (HIP) and consolidation of atomized powders. The HIP step was normally followed by hot forging of the hipped billets. This method of production is still the dominating powder metallurgical method to produce high speed steel.

The original objective for research and development on powder metallurgical processing of high speed steel was to improve its functional properties and performance in demanding applications. The main advantages from the powder metallurgical manufacturing process are no segregation and uniform and isotropic microstructure. The well known problems with coarse and severe carbide segregation in conventional cast steel and forged steel are thus avoided in powder metallurgical high speed steel.

Thus, the powder metallurgical manufacturing method of a high speed steel with sufficient amount of carbon and carbide forming elements, results in a disperse distribution of carbides that to a large extent solves the problem of low strength and toughness associated with conventionally produced high speed steel.

Hence, the present invention refers to a method for producing a high speed steel. The inventive method comprises the step of providing a powder consisting of the elements: 1-3 wt-% carbon (C), 3-6 wt-% chromium (Cr), 0-7 wt-% molybdenum (Mo), 0-15 wt-% tungsten (W), 3-14 wt-% vanadium (V), 0-10 wt-% cobalt (Co), 0-3 wt-% niobium (Nb),0-0.5 wt-% nitrogen (N), 0.2-1 wt-% yttrium (Y), and remainder iron (Fe) and unavoidable impurities, wherein Mo+0.5W =2-10 weight%. It should be pointed out that the elements having a lower limit of 0% are optional.

In an embodiment of the invention, the provision of the powder mixture comprises the step of argon gas-atomisation of molten metal comprising said elements into said powder. In a preferred embodiment of the invention, the argon gas-atomisation of the molten high speed steel causes high speed steel particles of a maximum size of 160 μm to be formed.

After the provision of the powder a body is formed from said powder. This forming may, for example, comprise pouring said powder into a capsule. The capsule is then evacuated, e.g. by being subjected to a negative pressure of below 0.004 mbar for 24 hours in order to evacuate said capsule. The capsule is then sealed in order to maintain said negative pressure in the capsule. The consolidation of the powder is achieved by subjecting the capsule to an elevated temperature, e.g. about 1150° C., and an elevated pressure, e.g. about 1000 bar, for a long period of time, e.g. two hours. This last consolidation step is called hot isostatic pressing, HIP.

A soft annealing step follows the HIP step, preferably the soft annealing step is performed at 900° C. followed by a temperature decrease to 700° C. at a cooling rate of 10° C./hour, from thereon the body is allowed to naturally cool down to room temperature.

After soft annealing the body may be subjected to machining and preferably a hardening (austenizing) step at 1100° C. and three subsequent annealing steps at 560° C. for 60 minutes each, with natural cooling to room temperature there between.

The resulting material from these subsequent steps exhibits a very good uniformity without the aforementioned segregations and coarse carbide structure, and the most important effect is that the yttrium element is evenly distributed in the base-matrix of the high speed steel.

TABLE 1 Molyb- Carbon Chromium denum Vanadium Tungsten Yttrium (C) (Cr) (Mo) (V) (W) (Y) Alloy wt-% wt-% wt-% wt-% wt-% wt-% A 1.28 4.2 5 3.1 6.4 — B 1.18 4.2 5 3.1 6.4 0.5 C 1.19 4.2 5 3.1 6.4 1 D 1.55 4 3.5 12 0.5 E 1.05 4 4.5 3.5 0.5

In order to demonstrate the superior properties of the inventive method, a high speed steel was designed without the optional elements, see table 1. The exclusion of the optional elements gives a clear and concise demonstration of the improved high-temperature wear due to the method. A simple evaluation method “pin-on-disc” for high-temperature wear is described below.

Table 1 shows the elements of the high speed steel used in the experiment. Smelts were produced with the elements in table 1, and from these smelts powders were produced by means of gas atomisation using argon. The powders of alloy B and C in table 1 have a particle size of <160 μm, while the powder of alloy A has a particle size of <500 μm.

In the following description, in order to further illustrate the present invention, a performed non-limiting experiment will be described in detail.

The preparation of samples continued with a filling of the capsules with powder, said capsules were made from spiral welded tubes with a diameter of 73 mm. The capsules were then exposed to an pressure below 0.004 mbar for 24 hours; the capsules were then sealed in order to maintain said pressure.

In order to consolidate the powder in the capsules a hot isostatic pressing operation was performed at 1150° C. and 1000 bar for 2 hours. The samples were then subjected to a soft annealing step at 900° C. followed by a temperature decrease to 700° C. at a cooling rate of 10° C./hour, from thereon the samples were allowed to naturally cool down to room temperature.

The samples were then machined and heat treated with a hardening (austenizing) step at 1100° C. and three subsequent annealing steps at 560° C. for 60 minutes each, with natural cooling to room temperature there between.

The final preparation step comprises stepwise grinding and polishing of the sample in automatic grinding/polishing equipment. During the final polishing step a 1 μm diamond suspension was used.

FIG. 1 shows a simplified test set-up used for the tribological testing; this set-up is known in the art and has been referred to as “pin on disc”. The principle for the “pin on disc” tribological testing is as follows; a sample 1 is rotated around an axis 5 with a speed ω for a number of revolutions. Simultaneous with the rotation of the sample 1, a force F is applied to a pin 2 that in its turn applies the same force F to a ball 3. The ball 3 is made of Al₂O₃ and has a diameter of 6 mm. The rotation of the sample 1 and the force F on the ball 3 causes a groove 6 to be formed in the sample 1.

In order to evaluate the wear behaviour at elevated temperatures, the lower part of the “pin on disc” set-up is accommodated in a furnace 4. Thus, the furnace 4 can heat the sample 1, the ball 3 and the lower part of the pin 2 to the desired operating temperature.

FIG. 2 shows a cross section of the groove 6 perpendicular to the longitudinal direction of the groove 6. The depth d, measured from the polished surface of the sample to the bottom of the groove 6, is used as a measure of the wear resistance of the sample. Another figure of the wear resistance is the cross-sectional area 7, which is defined as the cross-sectional area of the groove 6 below the polished surface of the sample 1 perpendicular to the longitudinal direction of the groove 6. The profile and depth d of the groove 6 was estimated using a Veeco Wyko NT9100 white light interferometer.

A series of samples according to the description above were produced and tested according to the “pin on disc” procedure outlined above. The “pin on disc” result is presented in FIG. 3, the linear speed in this test was 20 cm/s, the applied force F was 5N and 20N, respectively, and the samples were rotated 20000 revolutions.

As can be seen in FIG. 3, the addition of yttrium causes the depth of the groove to decrease at 650° C.; see alloy A with a groove depth d equal to 5.7 μm, alloy B with a groove depth d equal to 1.9 μm and alloy C with a groove depth d equal to 3.7 μm. This indicates the anticipated increased wear resistance at elevated temperatures for alloys produced by the inventive method. The addition of 0.5% yttrium to the high speed steel (Alloy B) caused a reduction of the groove depth d of roughly three times compared to the high speed steel without yttrium (Alloy A). Also the addition of 1% yttrium to the high speed steel (Alloy C) caused a reduction of the groove depth d at 650° C.

A more representative measure of the wear resistance is the volume loss per meter (mm³/m). The calculation of the volume loss per meter is performed by integrating the cross sectional area 7 over the longitudinal direction of the track and divide by the circumference of the groove. In FIG. 4 the volume loss per meter is presented; volume loss for alloy A is 4.6×10⁻⁵ mm³/m, volume loss for alloy B is 1.8×10⁻⁵ mm³/m and finally the volume loss for alloy C is 4×10⁻⁵ mm³/m. The relation between the yttrium content of the high speed steel and the volume loss per meter thereof is illustrated in FIG. 4. From FIG. 4, one can conclude that the yttrium content of 0.5% clearly results in the lowest volume loss per meter. A higher yttrium content than 1% also has a beneficial effect on the volume loss per meter. This relation implies that the yttrium content of 0.5% gives a superior increase in the implied wear resistance of the high speed steel. It should be noted that examples D and E, though not represented in the figures, also show corresponding positive effects due to the addition of yttrium thereto.

According to the invention the yttrium content of the high speed steel is within the range 0.2 to 1 weight %. It is preferred that the yttrium content of the high speed steel is more than 0.4 weight %, and less than 0.7 weight %, more preferably less than 0.6 weight %, more preferably 0.4 to 0.6 weight %, such as 0.4 to 0.5 weight %, such as 0.4, 0.41, 0.42, 0.43, 0.44, 0.45, 0.46, 0.47, 0.48, 0.49 and 0.5.

In FIG. 5, the hardness of the samples is presented. The hardness is 63 HRC for alloy A, the hardness is 57 HRC for alloy B and the hardness is 56 HRC for alloy C. The conclusion from FIG. 5 is that the hardness is reduced with the addition of yttrium. Without wishing to be bound to any specific theory, one possible explanation for this reduction is that less carbon is available in the alloys that contain yttrium, thereby reducing the hardness. This illustrates the theory that the wear rate of the high speed steel, in FIG. 3, at room temperature is primarily dominated by the hardness of the high speed steel. At room temperature the wear rate increases with decreasing hardness. However, at elevated temperatures other mechanisms are dominating the wear, such as the growth kinetics and the mechanical properties of the oxide scale. 

1. A method for producing a high speed steel that with reference to its chemical composition consists of the following elements, in weight %: 1-3 Carbon (C) 3-6 Chromium (Cr) 0-7 Molybdenum (Mo)  0-15 Tungsten (W)  3-14 Vanadium (V)  0-10 Cobalt (Co) 0-3 Niobium (Nb)  0-0.5 Nitrogen (N) 0.2-1  Yttrium (Y), and

the remainder iron (Fe) and unavoidable impurities, wherein Mo+0.5W=2-10 weight %, the method comprising the steps of: providing a powder comprising the elements of said high speed steel, forming a body of said powder; and subjecting said body to elevated heat and elevated pressure such that a consolidation of the powder thereof is achieved.
 2. A method for producing a high speed steel according to claim 1, wherein the step of providing the powder includes the step of argon-atomisation of molten metal comprising said elements into said powder.
 3. A method for producing a high speed steel according to claim 1, wherein the yttrium (Y) content of said high speed steel is more than 0.4 weight %.
 4. A method for producing a high speed steel according to claim 1, wherein the yttrium (Y) content of said high speed steel is 0.7 weight % or less.
 5. A method for producing a high speed steel according to claim 1, wherein the yttrium (Y) content of said high speed steel is within the range of from 0.45-0.60 weight %.
 6. A method for producing a high speed steel according to claim 1, wherein the carbon (C) content of said high speed steel is in the range of from 1.1-1.4 weight %.
 7. A method for producing a high speed steel according to claim 1, wherein the chromium (Cr) content of said high speed steel is in the range of from 3.0-6.0 weight %.
 8. A method for producing a high speed steel according to claim 1, wherein the chromium (Cr) content of said high speed steel is in the range of from 4.0-5.0 weight %.
 9. A method for producing a high speed steel according to claim 1, wherein the Molybdenum (Mo) content of said high speed steel is in the range of from 4.5-5.5 weight %.
 10. A method for producing a high speed steel according to claim 1, wherein the tungsten (W) content of said high speed steel is in the range of from 6-7 weight %.
 11. A method for producing a high speed steel according to claim 1, wherein Mo+0.5W=5.0-8.5 weight %.
 12. A method for producing a high speed steel according to claim 1, wherein the Vanadium (V) content of said high speed steel is in the range of from 3.0-5.0 weight %.
 13. A method for producing a high speed steel according to claim 1, wherein the Cobalt (Co) content of said high speed steel is in the range of from 8.0-9.0 weight %.
 14. A method for producing a high speed steel according to claim 1, wherein the high speed steel has a mean carbide particle size <3 μm.
 15. A method for producing a high speed steel according to claim 1, wherein the high speed steel has an isotropic microstructure. 